Various terms are defined in the following specification. For convenience, a Glossary of Terms is provided herein, immediately preceding the claims. As used herein, the terms “bead”, “weld bead”, “pass” and “weld pass” are synonymous, all of which terms are familiar to those skilled in the art of welding engineering. When used as a noun herein, the term “weld” means the same as “weld joint”, as will be familiar to those skilled in the art of welding engineering. A Table of References follows the Glossary of Terms. All REF. numbers referred to herein are identified in the Table of References. All publications listed in the Table of References are hereby incorporated herein by reference.
Steel is a widely used structural material in a variety of industries because of its low cost, desirable physical properties, and versatility through alloying, thermomechanical processing, and heat treatment. Within the past 10 to 15 years, significant advancements in steel making have enabled improved combinations of strength and toughness. Despite these advancements and continued broad use, one pitfall in the use of steel persists, the potential for low toughness and structural failure.
There are three primary reasons that the risk of low toughness persists. First, end users have responded to steel making improvements by selecting new, improved steels for more severe service conditions, e.g., applications that are less tolerant of oversights or unfortunate mishaps. Second, structural steel is a material with a body centered cubic crystal structure, therefore it displays a change from ductile behavior to brittle behavior as temperature is decreased. This aspect can be quantified by defining a ductile-to-brittle transition temperature (DBTT) which is a mathematical description of the position of the transition region on the temperature axis. FIG. 1 (Prior Art) shows a toughness transition curve 1 and the DBTT 2 is shown as the boundary between the upper transition region 3 and lower transition region 4. FIG. 1 has ordinate 5 representing increasing toughness and abscissa 6 representing increasing temperature. Lower shelf toughness region 7 and upper shelf toughness region 6 are also shown. Region 9 where cleavage (brittle) fracture may occur, region 10 where mixed mode fracture may occur, and region 11 where ductile fracture may occur, are also shown. The third reason that low toughness remains an issue is that the primary method used to join structural steel is fusion welding; and weldments, as compared to the highly processed base metal, often contain defects and degraded microstructures. Of particular concern is the toughness of the heat-affected zone (HAZ), a region adjacent to the fused weld metal where the base metal has experienced microstructural changes due to the heat from welding.
The coarse grain regions in certain areas of a multipass weld HAZ can be responsible for low toughness properties. As is explained below, these regions can be small and surrounded by material with relatively good toughness. A term used to describe small HAZ regions with low toughness is local brittle zones (LBZs). The topic of HAZ toughness in structural steels has received considerable attention in the literature, for example in REF. numbers 1 through 5. The annual Offshore Mechanics and Arctic Engineering conference proceedings is another example of literature on this topic, e.g., REF. numbers 3, 7, 10, and 12.
The Nature of Steel Heat-Affected Zones: Single Pass Welds
The microstructural formation of steel HAZs will be explained with an emphasis on material toughness. Mechanisms of toughness degradation due to metallurgical changes in steel HAZs fall into two general categories; (1) changes that hinder slip deformation, (2) changes that provide more potent cleavage initiation sites. Additional discussion about cleavage mechanisms in steel can be found in REF. number 5.
Single pass HAZs in steel can be divided into four regions based on the peak temperature reached during welding and its relation to the iron-carbon phase diagram. FIG. 2 (Prior Art) shows a hypothetical example of the four regions of a single pass HAZ for a steel with 0.10 wt % carbon content. The subcritical HAZ (SCHAZ) 15 undergoes only subtle changes due to its relatively low peak temperature. The high temperature boundary 12 of the SCHAZ 15 is the A1 transformation temperature and, therefore, the gross microstructure of the SCHAZ 15 remains untransformed. The low temperature boundary 14 is somewhat arbitrary and is not associated with any specific detail of the iron-carbon phase diagram. Low temperature boundary 14 is generally considered to be a temperature below which no significant changes occur to the original base metal. The definition of “significant” depends on the structural application. Metallurgical changes in the SCHAZ 15 are typically related to carbon/nitrogen diffusion, precipitation phenomena, and/or dislocation movement. These changes can be beneficial, benign, or detrimental to toughness depending on the specific steel and HAZ thermal cycle.
The intercritical HAZ (ICHAZ) 16 has the A1 transformation temperature and the A3 transformation temperature as its lower temperature boundary 12 and upper temperature boundary 18, respectively. In the ICHAZ 16, austenite will nucleate and grow during the higher temperature portions of the weld thermal cycle. The amount of austenite formation is small in regions near lower temperature boundary 12 of the ICHAZ 16 and is large (consuming most, if not all, of the microstructure) near upper temperature boundary 18. The properties of the ICHAZ 16 depend greatly on what the austenite transforms to upon cooling. The parts of the ICHAZ 16 that do not undergo much transformation to austenite can still experience significant diffusion of carbon, nitrogen, other alloys, and dislocation movement. The phase diagram shown in FIG. 2 has ordinate 13 representing temperature in ° C. and abscissa 17 representing weight percent carbon. Ferrite plus cementite region 22, alpha ferrite region 28, ferrite plus austenite region 19, austenite plus cementite region 27, austenite region 25, liquid plus austenite region 23, delta ferrite region 29, and liquid region 21 are also shown.
One mechanism of reduced toughness in the ICHAZ 16 is when the newly formed austenite transforms to a hard/brittle constituent on cooling. See, e.g., REF. numbers 5 through 7. FIGS. 3A-3F are from REF. number 5 and they provide schematics that describe this phenomena in a hypothetical ferrite-pearlite steel of about 0.10 wt % carbon. FIG. 3A shows a HAZ thermal cycle plotted on a graph having ordinate 30 representing temperature in ° C. and abscissa 34 representing time in seconds. The microstructure that exists at each of the three points indicated is shown in FIG. 3C-FIG. 3F. FIG. 3B shows a section of the iron-carbon phase diagram to explain the chemistry and microstructural changes in this HAZ region. The microstructure that exists at the point at about ordinate 0, abscissa 0, i.e., prior to welding, is a mixture of ferrite 31 and pearlite 32. During welding, once the temperature rises above the A1 transformation temperature, small “islands” of austenite 33 form within the ferrite 31′ and pearlite 32′. These islands of austenite 33 are enriched in carbon well beyond that of the base metal; e.g., islands of austenite 33 have about the carbon content indicated by point 38 in FIG. 3B, which is higher than the about 0.10 wt % carbon of the base metal. FIG. 3B has ordinate 35 representing temperature in ° C. and abscissa 36 representing weight percent carbon. Austenite region 39, ferrite plus austenite region 40, and ferrite plus cementite region 41, are also shown. If the cooling rate is sufficiently fast, the austenite may transform to martensite 37 or martensite-austenite constituent (M-A) 37 and the martensite can have a twinned substructure. This constituent can cause a reduction in toughness.
Referring again to FIG. 2, the fine grain HAZ (FGHAZ) 20 and the coarse grain HAZ (CGHAZ) 24 are adjacent areas that have been heated above the A3 transformation temperature, but below the steel's melting point. In general, the material within the FGHAZ 20 and CGHAZ 24 will transform completely to austenite during the welding thermal cycle, however, for some local areas this is not always the case. Due to chemical and microstructural inhomogenieties, the A3 transformation temperature can vary from point to point in the HAZ. It is possible that some local areas, or maybe just a few isolated grains, will not transform to austenite during the thermal cycle. Depending on the distance from weld fusion line 26, the austenite will grow to varying sizes depending mostly on peak temperature. Lower peak temperatures and smaller grains exist closer to the A3 transformation temperature boundary, and higher temperatures and larger grain sizes exist near the weld fusion line. The distinction between the FGHAZ 20 and CGHAZ 24 is arbitrary because, in reality, there is a continuum of grain sizes between the A3 transformation temperature boundary and the weld fusion line. Upon cooling, the austenite can transform to a number of different microstructures depending on the steel's chemistry and the cooling rate.
The terms “fine grain” or “coarse grain” for the FGHAZ 20 and CGHAZ 24, respectively, refer to the austenite grain size that existed during welding when the temperature was above the A3 transformation temperature. After cooling to ambient temperature, the austenite grains no longer exist, but there is, typically, a prior-austenite grain structure that is observable in an optical microscope. The prior-austenite grain size can have significant effects on HAZ toughness as explained in REF. number 5. The smallest prior-austenite grain size in the FGHAZ 20 can be on the order of a few microns whereas the largest prior-austenite grains in the CGHAZ 24 can be as large as 100 or 200 microns. Low heat input welds (e.g., about 1 kJ/mm) might display sizes of 75 to 100 microns near the weld fusion line. An approximate prior-austenite grain size to delineate the FGHAZ 20 from the CGHAZ 24 is on the order of 50 microns, plus or minus 20 microns, depending on the base metal grain size and the application. In other words, the FGHAZ 20 will contain a range of prior-austenite grain sizes from a few microns up to about 50 microns. The CGHAZ 24 will contain grain sizes from about 50 microns up to the sizes that exist near weld fusion line 26.
The properties of the FGHAZ 20 are dominated by the small grain size and the toughness tends to be very good. Near the low temperature boundary 18 of the FGHAZ 20, where the grain size is smallest, the toughness is often better than in the base metal. In contrast, the CGHAZ 24 is typically the lowest toughness region in a weld. A major factor is the relatively large prior-austenite grain size in the CGHAZ 24, but the final microstructure is significant as well. A range of microstructures can be produced in the CGHAZ 24.
The entire HAZ, from weld fusion line 26 to the outer boundary 14 of the SCHAZ 15, varies in width depending on the weld thermal cycle. For low heat input welds (e.g., less than about 1 kJ/mm) that tend to cool fast, the HAZ may be only a few millimeters wide. For higher heat input welds (e.g., around 3 to 5 kJ/mm) that cool more slowly, the HAZ may be about a centimeter wide. Similarly, the width of any single region, like the CGHAZ 24 or the ICHAZ 16, will be approximately a millimeter, or less, for a low heat input weld up to about a couple of millimeters, for a high heat input weld. When observing a HAZ that has been etched with a common chemical like nitol, only the CGHAZ 24, FGHAZ 20, and ICHAZ 16 show distinct etching. The SCHAZ 15 does not respond to such etchants because its gross microstructure is essentially unchanged (does not transform to austenite).
The Nature of Steel Heat Affected Zones: Multipass Welds
Most weldments in structural steel are multipass where successive beads are deposited one on top of the other. Each bead produces a HAZ that overlaps or crisscrosses some part of the HAZ of the previous pass. FIG. 4A through FIG. 4B (PRIOR ART) provide a schematic of the various regions in a two pass weld 42, including a first weld pass 43 and a second weld pass 44. The purpose of FIGS. 4A and 4B is to identify various HAZ regions including whether or not these regions have been altered by the heat from a subsequent weld pass. Regions 46 and 47 show unaltered HAZs from the first and second weld passes 43 and 44, respectively. Region 49 shows where a portion of the HAZ from the first weld pass 43 existed, but was eliminated by second weld pass 44. Region 48 shows the HAZ from the first weld pass 43 that was altered by the HAZ from the second weld pass 44. FIG. 4B shows that altered HAZ region 48 includes the intercritically reheated CGHAZ (IRCG) 50 and the subcritically reheated CGHAZ (SRCG) 51. FIG. 4B shows some unaltered areas adjacent to altered region 48, including unaltered CGHAZ 52, unaltered CGHAZ 53, unaltered FGHAZ 54, unaltered ICHAZ 55, and unaltered SCHAZ 56. The A1 transformation temperature isotherm 57 and the A3 transformation temperature isotherm 58 are also shown in FIG. 4B.
FIG. 5A through FIG. 5C (PRIOR ART) is a schematic of the various regions in a multipass weld. These regions include columnar weld metal 60, etched HAZ 61 in the weld metal, including CGHAZ, FGHAZ, and ICHAZ and also shows SCHAZ 62 in the weld metal. In somewhat greater detail, FIG. 5C shows unaltered CGHAZ 63, unaltered FGHAZ 64, unaltered ICHAZ 65, unaltered SCHAZ 66, SRCG 67, and IRCG 68. As shown by the markers 69 and 69′, the portion of the HAZ that was heated above the A1 transformation temperature can be made visible by etching. FIG. 4A through FIG. 4B and FIG. 5A through FIG. 5C are resketched from the publication, “Recommended Practice for Preproduction Qualification for Steel Plates for Offshore Structures”, API RP 2Z, Third Edition, August, 1998. The geometry of the weld shown in FIG. 5A through FIG. 5C is a half-K bevel where one plate edge is left unbeveled. The purpose of this weld geometry is to provide a “straight” HAZ for toughness testing. From a metallurgical standpoint, however, this schematic can be used to highlight the basic principles of multipass HAZ formation in any steel weldment.
FIG. 5A through FIG. 5C show that a multipass weld can produce HAZ regions that are either unaffected (e.g. unaltered areas) or significantly affected by the thermal cycles from subsequent passes. From the standpoint of fracture toughness, the principles previously discussed for single pass welds apply to the unaltered regions of a multipass weld and, usually, apply to subcritically reheated areas. Subcritical reheating generates a relatively low peak temperature, and the microstructure is generally unchanged from the original weld pass. Significant changes due to multipass reheating are typically associated with peak temperatures above the A1 transformation temperature.
The significantly altered regions in a multipass HAZ begin as one of the four single pass regions, and then upon the application of subsequent passes, they experience additional thermal cycles that change the microstructure. With respect to fracture toughness, one region deserves specific attention, the intercritically reheated CGHAZ (IRCG). “A Study Concerning the Heat Affected Zone Toughness of Microalloyed Steels,” PhD. Dissertation, D. P. Fairchild, The Ohio State University, Columbus, Ohio, June 1995, provides a schematic illustration of the microstructure that can form in the IRCG. This schematic is shown in FIG. 6A through FIG. 6G. FIG. 6A has ordinate 78 representing temperature in ° C. and abscissa 79 representing time in seconds. In this example, the base metal comprises about 0.10 wt % carbon and the beginning microstructure consists of ferrite 81 and pearlite 82. The first weld pass creates a coarse austenitic structure 85 at peak temperature and FIG. 6B shows that the coarse austenite is of the same carbon content 78 as it was at room temperature. FIG. 6B has ordinate 70 representing temperature in ° C. and abscissa 71 representing weight percent carbon. Liquid region 72, delta ferrite region 73, austenite region 74, ferrite plus austenite region 76, and ferrite plus cementite region 77, are also shown. Upon cooling most of the austenite 85 transforms to upper bainite 86. A small amount of proeutectoid ferrite 83 can also form from the austenite 85. During the second weld pass thermal cycle, once the temperature rises above the A1 transformation temperature, small “islands” of austenite 87 form, primarily, at prior-austenite grain boundaries. Some islands of austenite 87 also form on lath boundaries. The austenite islands 87 are enriched in carbon well beyond that of the base metal; e.g., austenite islands 87 have about the carbon content at point 75 in FIG. 6B, which is higher than the about 0.10 wt % carbon of the base metal. If the cooling rate is fast enough, the austenite 87 may transform to martensite 88 or martensite-austenite constituent (M-A) 88 and the martensite can have a twinned substructure.
The distribution of M-A islands 88 in the IRCG is somewhat different than in the ICHAZ. In the IRCG, the M-A islands 88 typically outline the prior-austenite grain boundaries. This morphology has been referred to as a “necklace” structure. Referring again to FIG. 3F, in the ICHAZ, the islands 37 are mostly located within regions that were base metal pearlite before welding and a necklace structure is not produced. Toughness degradation due to M-A islands in the IRCG can be severe. It has been the subject of numerous studies. For example, see REF. numbers 8 through 11. The primary effect of these islands is to create preferential sites for cleavage fracture initiation. This can cause a significant shift (deterioration) in the DBTT to higher temperatures.
The example and schematic provided in FIG. 6A through FIG. 6G is of a specific base metal and IRCG microstructure, but other possibilities exist. For example, it is possible that martensite, instead of upper bainite, will dominate the interior of the prior-austenite grains. It is also possible that the austenite islands on the prior-austenite grain boundaries will transform, on cooling, to a microstructure that is less detrimental than M-A. Regardless of the specific IRCG microstructure, two aspects are nearly unavoidable; (1) enlarged prior-austenite grains and, (2) a necklace arrangement of small grains outlining the prior-austenite grains. While some IRCG improvements can be produced by altering the base metal chemistry or welding procedure, the large grain size and necklace structure are characteristic, and they cause toughness reduction in the IRCG relative to other HAZ regions and/or the base metal.
HAZ Toughness and Structural Integrity
With respect to the integrity of structural weldments, the primary mechanical properties of interest are strength and toughness. Strength provides general load carrying capability while toughness provides load carrying capability when defects are present. As is known to those skilled in the art of welding engineering and structural mechanics, weld defects are a fact of life, and these flaws can cause failure by brittle fracture at loads below the design criteria.
When loads are applied to a steel component that contains a sharp defect (like a crack), a stress concentration is created near the defect tip. The elevated stresses may only affect a region a few millimeters in diameter. Despite this region being small, if the local microstructure lacks sufficient toughness, then a cleavage fracture can initiate. In fact, it is possible for cleavage fracture, and subsequent structural failure, to initiate within a region that is just several microns in cross sectional diameter.
The local nature of cleavage initiation provides a mechanism whereby small, low toughness HAZ regions (i.e., LBZs) can cause structural failure even if these regions are surrounded by tough material. The possibility of LBZs causing structural failure has been debated primarily within the engineering community that designs, produces steel for, and/or builds fixed offshore platforms. For example, see REF. numbers 1 through 4, 6, 7, and 10 through 12. After about twenty years of study, it is generally believed that the mechanism of cleavage initiation from LBZs is real. It is not a testing fluke associated with laboratory measurements. On the other hand, the absence of platform failures due to LBZs requires an explanation. It appears that while this failure mechanism is possible, there are enough safety factors in place (for offshore platforms) that the probability of failure is low. In other words, the probability is low that a very sharp defect will be located in or near an LBZ in an area of a structure that will experience the simultaneous occurrence of a large enough load and a cold enough temperature to cause failure.
For applications other than offshore platforms, determining the significance of LBZs requires attention paid to topics like material toughness, loading type (fatigue, static, impact, etc. . . . ), structural redundancy, in-service inspection, and fabrication methods. While some applications will be inherently resilient to LBZs, it is likely that others will prove sensitive to LBZs and new techniques will be necessary to provide suitable designs. Considering the trend of ongoing steel making improvements and the selection of these materials for harsh service, the development of welding methods that increase the toughness of LBZ regions, or eliminate them entirely, would be very useful.
U.S. Pat. No. 1,554,546
It has been known for some time that the multiple thermal cycles imposed on any one HAZ region in a multipass weld can be beneficial for certain mechanical properties. In fact, welding procedures have been specifically designed to take advantage of multipass heating for the improvement of properties like HAZ toughness, corrosion resistance, and hydrogen cracking resistance. See, for example, REF. numbers 13 and 14. Related to this type of procedure, J. B. Austin, in U.S. Pat. No. 1,554,546, describes a method where a weld bead, called a “refiner”, is placed on top of another for the purpose of improving the properties. Austin describes that some welds contain undesirable factors such as “internal stresses” and coarse, brittle microstructures. He states that his refiner weld bead can reduce or eliminate these factors. Several of the drawings from Austin's patent are redrawn in FIG. 7 through FIG. 12A. FIG. 7 and FIG. 8 show schematic before and after cross sections of a weld 90 with the refiner bead 91 applied. FIG. 9 and FIGS. 10 and 12A (copy of FIG. 10) show fillet and butt weld geometries, respectively, that have refiner beads 91′ and 91″ applied. FIG. 11 shows a fillet weld with two refiners 93 and 94 applied.
In discussing improvements to welds, or as Austin refers to them “seams”, Austin states, e.g., at page 2, lines 122-130, that by the application of his invention, one will have “ . . . prevented the formation of or removed any embrittlement in the base metal or metals adjacent to the line of fusion of such seams.” Austin states, e.g., at page 3, lines 126-130, that his refiner weld bead “apparently . . . reheats the seam weld and zones of the base metal . . . up to or somewhat above the A3 temperature . . . ” and he states, e.g., at page 4, lines 24-30, that the “ . . . adjacent base metal is grain refined and further, any brittle constituents of the base metal adjacent to the weld are substantially eliminated.”
J. B. Austin's patent is dated 1925, and at that time, relatively little was known about the details of metallurgical transformations in steel HAZs. In fact, the term “heat affected zone” was not yet in use. The definitions and detail shown in FIG. 1 through FIG. 6G, of this application, and the knowledge of how HAZs change with steel chemistry were unknown to Austin. It was fortuitous that J. B. Austin noticed some beneficial effects of his “refiner” weld, because by present day standards, it can be anticipated that certain regions within his welds were either not improved over the original weld or they may have been degraded. If Austin's butt weld, as shown in FIG. 10 (of this application), is redrawn to show various HAZ regions, then FIG. 12A and FIG. 12B (of this application), showing the base plate 98 and refiner weld 99, results. In FIG. 12A and FIG. 12B, the presence of unaltered CGHAZ 100, IRCG 102, and SRCG 104 regions are shown. As discussed above, these regions typically have low toughness. Certainly the unaltered HAZ of the refiner bead would have no better toughness than the original unaltered HAZ of the primary bead. It can be reasoned that while some areas in J. B. Austin's welds were improved due to multipass reheating, there still existed areas of low toughness, i.e., LBZs.
In 1925, LBZs could not have been detected because no test methods were known for measuring the appropriate property, fracture toughness. The lack of metallurgical understanding of the time prevented the development of suitable test methods. Also, the significance of sharp defects in steels was not known (see, e.g., REF. number 5) and the idea that the integrity of a weld joint could be compromised by small low toughness regions was a foreign concept. In 1925, tests of ductility or “brittleness” did not incorporate the use of sharp cracks and the evolution of fracture mechanics was still decades away.
Another reason that it would have been difficult in 1925 for anyone to understand the shortcomings of J. B. Austin's patent is related to the quality of steels and welds. The idea of degraded toughness presumes that a comparison toughness is available that is relatively high. Because of the general low quality of steel and welds in 1925, a technique like Austin's was considered an improvement. However, if the same technique is applied at the time this patent application is filed, it might be considered detrimental. In a fortuitous sense, Austin uncovered some phenomena about the benefits of multipass welding, but in reality his methods have modern day limits because his techniques fail to account for metallurgical factors like LBZs and the structural significance of such regions.